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In-situ investigation of phase transformation behaviors of 300M steel in continuous cooling process
Materials Characterization 144 (2018) 400–410
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In-situ investigation of phase transformation behaviors of 300M steel in continuous cooling process
T
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Rongchuang Chen, Zhizhen Zheng , Ning Li, Jianjun Li, Fei Feng School of Materials Science and Engineering, Huazhong University of Science and Technology, and State Key Laboratory of Materials Processing and Die & Mould Technology, 430074, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Microstructure evolution Phase transition In-situ investigation 300M steel
Comprehensive understanding of continuous cooling transformation behaviors is important in heat treatment of steels, but there is still a lack of research in the phase transition behaviors of 300M steel, which would greatly hinder the microstructure controlling and finally affect the application of this material. In this work, in-situ observations by high temperature confocal laser scanning microscope (HTCLSM) was introduced by considering its superiorities in clarity, accuracy, and directness, to systemically investigate the microstructure evolutions of 300M steel under various cooling rates (0.01–100 °C/s). Pearlite, bainite, and martensite were observed to form at the cooling rate range of 0.01–0.15, 0.03–1, and 0.3–100 °C/s, respectively. A continuous cooling transformation diagram was constructed based on the in-situ observations, verified by metallography, and compared with dilatometry. Results showed that the transition temperatures by in-situ observations agreed well with the results of dilatometry, while the transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher. Finally, models accurately describing the relationships between Vickers hardness, retained austenite content, and cooling rates were established, and the phase transformation mechanisms and kinetics were analyzed. Our finding not only understands fundamentally the transformation mechanisms of different microstructures, but also provides a useful reference for practical microstructure control in heat treatment of 300M steel.
1. Introduction Steel, as a high strength and economic material, has played an extraordinary role in the development of modern industry. As a kind of ultra-high strength steel (σb ≥ 1800 MPa), medium alloyed 300M steel has been widely used in aircraft landing gear, airframe parts, and pressure vessels owing to its excellent mechanical properties [1]. However, 300M steel is facing difficulties in microstructure control in heat treatment, which seriously affects the mechanical properties and the safety of components in service. In general, the microstructures of steel could be theoretical analyzed and controlled through heat treatment, therefore the continuous cooling transformation (CCT) diagram or time-temperature-transformation (TTT) diagram becomes crucial for steels [2]. Compared with the TTT diagram, the CCT diagram has attracted more attention because it has more similarity to the cooling process in industrial applications. Although the CCT diagram of 300M steel was established by Dong et al. [3], very few in-situ observations of phase transformations of 300M steel were found, and the researches on the phase transformation behaviors of 300M steel are still needed.
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Various methods were adopted by researchers to gain a better insight into the phase transformation behaviors of steels. In conventional continuous cooling and interrupted cooling method, several groups specimen were separately cooled at constant cooling rates and quenched at different temperatures. The starting and terminating temperatures for phase transformations were determined by etching and metallography, and subsequently the CCT diagram was constructed [4]. It is worth noting that quenching might destroy the high temperature microstructures and therefore affect the accuracy of results [5]. Therefore, efforts have been made to solve this problem. One is dilatometry, in which a dilatometer recorded the thermal expansion as a function of temperatures, and then the phase transition temperatures could be determined [6,7]. Other methods such as ultrasonic techniques, high temperature X-ray diffraction (XRD), and differential scanning calorimeter analysis were also implemented to characterize the phase transformations of alloys [8–11], which significantly reduced the number of experiments. However, some factors such as thermal expansion, sound wave propagation velocity, and the width of the XRD profile were influenced by grain growth and dislocation density as well
Corresponding author. E-mail address: [email protected] (Z. Zheng).
https://doi.org/10.1016/j.matchar.2018.07.034 Received 3 May 2018; Received in revised form 14 July 2018; Accepted 26 July 2018 Available online 27 July 2018 1044-5803/ © 2018 Published by Elsevier Inc.
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as phase transitions, which inevitably introduced experimental error [12,13]. In addition, if two phases precipitated successively in cooling, it was very hard to distinguish between their phase transformation temperatures from the measured curves [6,14]. Furthermore, the relationships between temperatures and properties, such as the thermal expansion and specific heat capacity, were usually nonlinear even under no phase transformation conditions, which caused difficulty in transition temperature determinations [1]. Besides, the phase transformations were not directly observed by these methods, and the results by these methods should be further validated by metallographic observations. Accordingly, a more direct method is urgently needed. Comparing with these conventional methods, in-situ observation provided precise, direct, and continuous observations without altering the testing conditions. In general, high temperature optical microscope (HTOM), high temperature scanning electron microscope (HTSEM), and high temperature confocal laser scanning microscope (HTCLSM) are the most successful techniques for in-situ observations. But very few scanning electron microscopes could be adopted in high temperature observation of steels, which hindered the engineering applications in phase transformation researches [15]. Several in-situ observation works of phase transformations based on HTOM have been carried out, but difficulties were encountered when the phase transformation mechanisms needed to be explained in detail at higher magnifications (≥500) [16–18]. The confocal laser scanning microscope (CLSM) improved the optical imaging precision by using a spatial pinhole to block out-offocus light in image formation and has been successfully used in life sciences, pharmacology, and materials research, etc. The high temperature confocal laser scanning microscope (HTCLSM), an instrument developed for metal high-temperature metallographic observation based on CLSM, could simulate the heat treatment process of industrial conditions via accurate process controlling in protective atmospheres. The HTCLSM had its distinctive advantages in imaging clarity comparing with HTOM, and compared with HTSEM, the process control accuracy was higher, but the cost was much lower. Recently, several researches on alloys solidification, phase transformation, and grain growth have been reported [19–22] based on the HTCLSM. Although the phase transformation and grain growth of various alloys have been investigated, there is no in-situ research on the phase transformation of 300M steel. Accordingly, the HTCLSM was adopted to investigate the continuous cooling transformation behaviors of 300M steel. The detailed phase transformation behaviors at various cooling rates will be investigated. The CCT diagram of 300M steel will be established according to in-situ observation and compared with the result of dilatometry. This investigation intends to establish a detailed and systematic description of the phase transformation behaviors of 300M steel during continuous cooling.
2.2. XRF and SEM-EDS Tests The compositions were measured via an XRF apparatuses (XRF1800, Shimadzu Inc., Kyoto, Japan), and the error of element content measuring was claimed to be within ± 0.01% by the instrument supplier. The compositions in Table 1 showed that the segregation basically did not exist in the as-received ingot, although the carbon content of half radius was a little higher, which was within the allowable range of experimental error. The microstructures of as-received material and heat-treated material were observed on a scanning electron microscope (SEM, Quanta650 FEG, FEI Co., America) and the chemical compositions were quantitatively analyzed by cathodoluminescence spectroscopy (MP-32S, Horiba Co, Japan). 2.3. In-situ Observation Experiments The microstructure transformations during continuous cooling of 300M steel were in-situ observed by HTCLSM. The HTCLSM setup (VL2000DX-SVF17SP, Yonekura MFG Co., Japan), schematically illustrated in Fig. 1c, used a spatial pinhole to block out-of-focus light on a laser scanning microscope to achieve high resolution (0.14 μm) imaging. The Maximum scanning speed was 120 frames per second, ensuring that enough intermediate steps could be captured during phase transitions. In the case of quenching at 100 °C/s, the cooling time was ~11 s, and the maximum number of frames that could be captured was 1320, meaning that the HTCLSM could still record the martensitic transformation process by capturing one picture per degree centigrade decreasing. The temperature was measured using a thermocouple connected to the bottom of ceramic crucible on which the observed specimens were put. The temperature was continuously adjusted by a computer to obtain the specified temperatures and cooling rates. The temperature controlling error was claimed to be within ± 0.1 °C by the instrument supplier. During continuous cooling, the boundaries of different grains and phases at high temperatures could be visualized because of the volatilization of low melting point alloying elements at interfaces and the surface reliefs generated due to phase transformation. It was reported that the surface reliefs for bainite were inversed ‘V’ shapes and for martensite they were ‘N’ shapes [1,23]. The specimens were heated at 1 atm argon atmosphere to 1150 °C at a constant rate of 3.33 °C/s, held for 10 min, and then cooled to 25 °C at various rates ranging from 0.01, 0.03, 0.05, 0.1, 0.15, 0.3, 0.5, 1, 2, 3.33, 10, to 100 °C/s. The test specimens for HTCLSM were cut from ingot and turned to the size of 6 mm in diameter and 3 mm in height, and then mechanically polished according to standard metallographic procedure [24]. 2.4. XRD Tests
2. Materials and Experiments
The retained austenite contents of heat-treated specimens were quantitatively analyzed via an X-ray diffraction instrument (D8 Advance, Bruker Co., Germany). The specimens for XRD tests were square sheets with the size of 10 mm × 10 mm × 2 mm, and the measured surfaces were mechanically polished. The scanning speed was 1°/ min. The Co target was used in XRD experiments. The retained austenite contents were calculated via a spectrum analysis software (Topas, Bruker Co., Germany) automatically.
2.1. Materials The 300M steel ingot was in an annealed state with the diameter of 300 mm. The chemical compositions of the ingot were measured via XRay Fluorescence (XRF). In order to investigate the segregation behaviour of the ingot, specimens for XRF tests were sampled from the center, the half radius, and the outer edge of the ingot by wire-electrode cutting. Then the specimens were turned to 33 mm in diameter and 12 mm in height, and mechanically polished. The chemical compositions on different locations of the ingot were shown in Table 1. A scanning electron microscope integrated with energy dispersive spectrometer (SEM-EDS, Quanta650 FEG, FEI Co., America) was used to characterize both microstructure and chemical compositions, as shown in Fig. 1a and b. The microstructure was upper bainite. The chemical compositions measured by EDS (Fig. 1b) agreed well with the XRF result.
2.5. Microhardness Tests Micro-hardness of cooled specimens was tested on a Wilson Hardness tester (430SVD, Buehler Co., America). The test error was claimed to be within ± 0.75%, but the errors could be increased because the diagonal lengths of indentations were measured by visual inspection. The loading force was 9.8 N, and the loading time was 10 s. At least 5 points on each specimen were tested. 401
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Table 1 Chemical compositions of 300M steel (mass fraction %). Element
C
Mn
Si
Ni
Cr
V
Mo
S
Fe
Center Half radius Outer edge
0.39 0.40 0.39
0.801 0.810 0.814
2.707 2.362 2.618
1.822 1.833 1.818
0.897 0.901 0.891
0.085 0.087 0.085
0.432 0.435 0.437
0.019 0.016 0.017
Balanced Balanced Balanced
2.6. OM, SEM, and TEM Characterizations
3. Results and Discussion
Since the conventional etching and metallographic analysis were proved effective in most conditions, the cooled specimens at different cooling rates were characterized by etching and observation on OM and SEM. Specimens for OM observations were mechanically polished and etched using etchant which contained 4% HNO3 and 96% CH3CH2OH (volume fraction). An optical microscope (VHX-1000C, Keyence Co., Japan) was used for observations. Specimens were then photographed on a scanning electron microscope (SEM, Quanta650 FEG, FEI Co., America). The acceleration voltage was 10 kV. Samples for TEM tests were mechanically polished to the thickness of ~70 μm, and electro polished to obtain the thin areas for photography. The etchant for electro polishing contained 10% HClO4 and 90% CH3CH2OH. The operating voltage was 20 V, and the temperature was −20 °C. A transmission electron microscope (TECNAI G2 F30, FEI Co., America) was used in microstructure characterization. The acceleration voltage was 200 kV.
3.1. The Construction of CCT Diagram by In-situ Observations Fig. 2 showed the microstructure transformations with decreasing temperatures under various cooling rates. For example, under cooling rate of 0.01 °C/s, lamellar structures gradually formed in austenite grains with lamellar spacing decreased from 6.9 to 2.4 μm as temperature decreased from 840 (Fig. 2a) to 781 °C (Fig. 2b), and this tendency became conspicuous, such as the lamellar was almost unrecognizable on HTCLSM when the temperature further decreased to 709 °C (Fig. 2c), and the morphology of the lamellar structure remained unchanged thereafter. The interface marked in Fig. 2a–c, also named as growth ledge in literature [25], did not migrate during cooling. Under cooling rate of 0.15 °C/s, some part of the grains became darker from 753 (Fig. 2d) to 702 °C (Fig. 2e), which suggested the reflectivity changing owing to the microstructure transformations of the grains. It was also observed in Fig. 2f that upper bainite was generated when temperature decreased from 492 to 345 °C while the appearance of some grains remained unchanged throughout the cooling. Under cooling rate of 0.3 °C/s, no microstructure transformation could be seen until the temperature drops to 493 °C (Fig. 2g), which located in the temperature range of bainite transformation, surface reliefs were observed to form at an average speed of about 1.5 μm/s across grains and at grain boundaries. Fig. 2h and i showed the detailed transformation progress. Under the cooling rate of 3.33 °C/s, the microstructure transformation only occurred below 295 °C (Fig. 2j), wherein lath
2.7. Thermal Expansion Tests To compare with the phase transition points determined by HTCLSM observations, the thermal expansions were recorded by dilatometry. Specimens were turned to cylinders with the size of 8 mm in diameter and 12 mm in height. Thermal expansions were measured on a thermal compression machine (Gleeble 3500, Dynamic Systems Inc., America) at cooling rates of 0.15 and 0.3 °C/s, respectively.
Fig. 1. As-received material and experiment schematic. (a) SEM graph showed the microstructure was bainite; (b) EDS spectrum of as-received 300M steel; (c) schematic illustrations of the HTCLSM setup. 402
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Fig. 2. Microstructure transformations at cooling rates of (a)–(c) 0.01 °C/s; (d)–(f) 0.15 °C/s; (g)–(i) 0.3 °C/s; (j)–(l) 3.33 °C/s by in-situ observations on HTCLSM. “Lath A–F” in Fig. 2j–l denoted martensite laths. “GB” meant grain boundary. “UB” was upper bainite. The corresponding temperatures were marked below the photos. (×100).
divided into 7 regions, the microstructures of each were shown. The upper right region (A) showed the equiaxed austenite grains before transformations. At slow cooling rates, pearlite lamellas and proeutectoid ferrite (P + F) were the dominant structures. As the cooling rate increased, the microstructure transited from mixed microstructures of pearlite and bainite (P + B), bainite and martensite (B + M), to martensite (M). It should be noticed that the incomplete transformation was a common characteristic of bainite and martensite transformations, so there was retained austenite in the transition product as well, which should be further investigated. The regions marked ‘A → B’ and ‘A → M’ showed the intermediate state of bainite and martensite transformations. The pearlite, bainite, and martensite formed at the cooling rate range of 0.01–0.15, 0.03–1, and 0.3–100 °C/s, respectively. The transformation starting temperatures were within the range of 753–840,
martensite appeared very fast across grains and grew more and more as temperature decreased, as shown in Fig. 2k and l, respectively. It was interesting that the newly formed martensite laths were connected end to end, and almost all laths terminated at grain boundaries (marked in Fig. 2j) or where they met (Fig. 2k). The zoomed views of the pearlite formation process at the cooling rate of 0.01 °C/s were shown in Fig. 3. In the starting stage of pearlite formation, lamellas with relative big spaces were formed (Fig. 3a) in some austenite grains, and the zoomed views in Fig. 3b showed the oval nuclei between pearlite lamellas. With decreasing temperature, it could be seen from Fig. 3c and d that the newly formed pearlite lamellas densified the previously formed pearlite lamellas. To further investigate the effect of cooling rates on microstructure transformations, the HTCLSM results were summarized in Fig. 3. It was 403
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Fig. 3. Zoomed views of the pearlite formation process at the cooling rate of 0.01 °C/s and photographed at (a) 832 °C, (b) 832 °C, (c) 756 °C, and (d) 730 °C.
444–496, and 295–297 °C. At cooling rate of 0.03–0.15 (pearlite and bainite transformation) and 0.3–1 °C/s (bainite and martensite transformation), two transformations occurred successively. The XRD profiles of cooled specimens at various cooling rates were shown in Fig. 5a, according to which the retained austenite contents could be calculated, shown in Fig. 5b. The retained austenite was marked out using face centered γ-Fe, and the α-Fe represented body centered ferrite or martensite depending on the cooling rates. As could be seen, with increasing cooling rate, the intensity of γ-Fe peaks decreased gradually, indicating the retained austenite content decreased. Sherman et al. [26] reported an inverse liner relationship between the retained austenite contents and cooling rate of a martensitic steel in the cooling rate range of 25–560 °C/s. The result of Li et al. [27] also showed that the retained austenite content of 0.19C-1.5Si-1.6Mn steel decreased with increasing cooling rate from 0.05 to 10 °C/s. In the present investigation, in a wider range of cooling rates, the retained austenite content (Cra) was found to decrease with cooling rate (v), and the relationship between the retained austenite contents and the logarithm of cooling rate could be fitted using a second order polynomial:
Cra =
2 0.009496⋅log10 v
− 0.04329⋅log10 v + 0.04563
Fig. 4. Relationships between microstructure transformations and cooling rates. The letters “A”, “P”, “B” and “M” denoted austenite, pearlite, bainite, and martensite, respectively.
(1)
The symbols and line in Fig. 5b showed the experimental and fitted values. The R value in representative of the confidence level was calculated to be 0.937, indicating that the Eq. (1) was moderately precise in calculating the retained austenite content. The influence of cooling rates on retained austenite content could be explained by Fig. 4. At slower cooling rates, the transformations of γ-Fe to α-Fe were slower as well, resulting in a relative longer time to complete the transformations, and a higher content of retained austenite was obtained. Meanwhile, at slower cooling rates, the strains generated in phase transformations, which was beneficial for phase transformation according to Zhang [28], were more easily to be eliminated by recovery. Therefore, the retained austenite content was higher in the slow cooling. The micro-hardness of cooled specimens was tested, as shown in
Fig. 6. Basically, the relationships between the Vickers hardness and the logarithm of cooling rates followed the S-curve. The hardness increased with increasing cooling rates until at 3.33 °C/s the hardness became stable at HV 665 ± 15. It was confirmed by HTCLSM results that the main transformation product was pearlite, bainite, and martensite at the cooling rate range of 0.01–0.15, 0.03–1, 0.3–100 °C/s, respectively. Therefore, the bainite seemed to be an intermediate of pearlite and martensite in terms of hardness. The relationships between cooling rates (v) and Vickers hardness (Hv) could be fitted as:
Hv = 667.92 −
404
286.65 1 + exp(2.3013⋅log10 v + 1.4117)
(2)
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Fig. 5. Retained austenite content measured by XRD. (a) XRD profiles; (b) retained austenite content in percentage. Table 2 Phase transition points. The “M”, “B”, and “P” denoted martensite, bainite, and pearlite. The subscript s and f denoted phase transition starting and finishing temperatures, respectively. The symbol “-” denoted the absence of the corresponding transformations. v (°C/s)
0.01 0.03 0.05 0.1 0.15 0.3 0.5 1 2 3.33 10 100
Fig. 6. Hardness of specimens cooled at different cooling rates.
Phase transition temperatures (°C) Ms
Mf
Bs
Bf
Ps
Pf
– – – – – 296 295 295 295 295 299 297
– – – – – 160 160 164 169 160 160 162
– 496 492 496 492 493 460 444 – – – –
– 369 359 352 345 329 311 299 – – – –
840 821 800 770 753 – – – – – – –
709 708 706 708 702 – – – – – – –
began not to change. The transformation starting and finishing temperatures of martensite almost kept constant at 295 and 160 °C when the cooling rate varied from 100 to 0.3 °C/s, but the martensite transformation disappeared during cooling at a rate higher than 0.3 °C/s, so at that point, the Ms and Mf points were connected with a dotted line, indicating a termination of martensite transformation. With increasing cooling rates from 0.03 to 1 °C/s, the transformation starting and finishing temperatures of bainite decreased from 496 to 444 and from 369 to 299 °C, respectively. The transformation starting temperatures of pearlite also showed a significant decrease from 840 to 753 °C as the cooling rates increased from 0.01 to 0.15 °C/s while the finishing temperature almost kept constant at 706 °C by the HTCLSM result. It could be explained that for diffusion type transformations, for example, the bainite and pearlite transformations, the transformation temperatures would be lower at fast cooling rates due to insufficient diffusions. Besides, with decreasing cooling rates, the lower-bainite transformation and the upper-bainite transformation occurred sequentially, and the upper-bainite transformation temperatures were usually higher than the lower-bainite transformation temperatures.
Fig. 7. Continuous cooling transformation diagram of 300M steel.
The R value is 0.962, which meant the model could calculate the Vickers hardness of 300M steel with high precision in a wide cooling rate range from 0.01 to 100 °C/s. According to the in-situ observations, the CCT diagram was constructed and shown in Fig. 7, and the corresponding phase transition points were shown in Table 2. The starting points were estimated to be the first occurrence of newly generated microstructures, and the termination points were the temperatures at which the microstructure
3.2. Verification by Traditional Methods Since the conventional etching and metallographic analysis were proved effective, the cooled specimens at different cooling rates were 405
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Fig. 8. Etching results of 300M steel austenitizing at 1150 °C for 10 min and cooled at (a) 0.01; (b) 0.15; (c) 0.3; (d) 3.33 °C/s. PF denoted proeutectoid ferrite. BF meant bainite ferrite. (×200).
3.3. Comparison With the Results of Dilatometry
characterized by etching and then observed on OM. Fig. 8 showed the etching results of cooled specimens at various cooling rates. It showed that at the cooling rate of 0.01 °C/s (Fig. 8a), the equilibrium microstructures containing proeutectoid ferrite and pearlite with lamellar structures were formed. At 0.15 °C/s (Fig. 8b), the mixed microstructures of bainite ferrite and lamellar pearlite were formed. At 0.3 °C/s (Fig. 8c), lower bainite was dominant and some martensite laths were observed as well. At cooling rates of 3.33 °C/s, lath martensite was the main microstructure, as shown in Fig. 8d. The microstructures characterized by traditional methods agreed well with the result by HTCLSM. In order to scrutinize the surface morphology and substructure and get a better understanding of transformation mechanisms of different microstructures, cooled specimens obtained at various cooling rates were characterized by SEM and TEM and shown in Fig. 9. At 0.05 °C/s, mixed microstructures of pearlite and bainite were observed (Fig. 9a). The shape of pearlite was irregular lamellae, which indicated spheroidization of the cementite (Fig. 9b). It also showed that acicular bainites were generated and ferrites were distributed between bainites (Fig. 9c). At cooling rate of 0.15 °C/s, feathery bainites formed at grain boundaries (Fig. 9d). It was found that the most frequent angle between bainite leaves and grain boundaries were 43°. When cooled at 100 °C/s, martensite laths with the angle of 57–72° were generated (Fig. 9e). It could be seen from Fig. 9f that the martensite laths were blocked by the laths with angle of ~65°. The substructures of martensite lath were dislocations, as shown in the dark areas in Fig. 9g that the dislocation piled up in a martensite lath. The microstructures characterized by SEM and TEM agreed well with the result by HTCLSM.
In order to inspect the reliability of the CCT diagram established in by in-situ observations, the results were compared with the results of dilatometry. Thermal expansions were measured at the cooling rate of 0.15 and 0.3 °C/s and compared with the HTCLSM results, as shown in Fig. 10, where the phase transition temperatures obtained by HTCLSM observations were marked as well. Although the phase transition points could be observed by the thermal expansion method, it was very difficult to obtain accurate phase change points by dilatometry method especially in the stage when martensite and bainite were transformed successively. The transition points by in-situ observations agreed well with the results of dilatometry. But the transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher. It could be explained that some transformations might have taken place prior to the observed transformation points and not been observable because of a too low surface relief. The same reason was applicable to the explanation of the termination points of phase transitions. Although it may be less rigorous to declare that the result of in-situ observation was more accurate than the result of dilatometry, the in-situ observations by HTCLSM does have an obvious advantage in terms of directness and convenience.
3.4. Phase Transition Mechanisms Uneven carbon distribution existed in austenite grains before the transformation, which can be proved by the result in Fig. 2a that the
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Fig. 9. SEM and TEM results of 300M steel showing the various microstructures at different cooling rates: (a) 0.05 °C/s, P + B, SEM, ×750; (b) 0.05 °C/s, P, SEM, ×4000; (c) 0.05 °C/s, B, SEM, ×2000; (d) 0.15 °C/s, P + B, SEM, ×750; (e) 100 °C/s, M, SEM, ×2000; (f) 100 °C/s, M, TEM, ×3800, dark field image; (g) 100 °C/s, M, TEM, ×100,000, bright field image.
uneven carbon distribution was aggravated by the long-range carbon diffusion at a slow cooling rate, as shown in Fig. 8a that microstructures consisted of proeutectoid ferrite and spheroidal pearlite were formed at 0.01 °C/s. At moderate cooling rate, the proeutectoid ferrite was not formed because the long-range diffusion of carbon could not be completed due to lack of time. Therefore, the carbon was retained in ferrites, and some of the carbon precipitated from ferrite, forming cementite. This was proved by the results in Fig. 9c and d that the cementite distributed among bainite ferrite. The transition processes of pearlite, bainite, and martensite were illustrated in Fig. 11 according to the in-situ observations. As shown in Fig. 11a–d, at a slow cooling rate of 0.01–0.15 °C/s, as the solubility of carbon in supercooled austenite decreased with increasing supercooling degree, equidistant, parallel, and necklaced cementite nucleated in austenite grains. Then, the coarse cementite lamella with several microns spacing were initially formed, consuming carbon in super cooled austenite. As a result, the carbon content of the regions far from the cementite layers was relatively high comparing with the regions adjacent to the cementite layers, so the cementite nucleated and grew, leading to a reduction in lamellar spacing. Fig. 11e showed that cementite nucleated between cementite layers. It should be noticed that the pearlite lamella formed in a gradually denser way, as shown in Fig. 11a–d, and the newly formed lamella densified the lamella formed at higher temperatures. The so called “growth ledge” [25] in Fig. 2a–c did not migrate during the formation of pearlite lamella, which meant
Fig. 10. Comparison of the phase transition points obtained by in-situ observations and by dilatometry during cooling at 0.15 and 0.3 °C/s.
austenite grains did not transform with simultaneously. The higher carbon content, the bigger the resistance of phase transformation, and therefore the lower the phase transition temperatures would be. It was reported by Vermeulen et al. [29] that the Ms would decrease by about 275–375 °C for every 1% of equivalent carbon content decreasing. The 407
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Fig. 11. Illustration of the pearlite, bainite, and martensite formation. Figures in (a)–(d), (f)–(i), and (k)–(n) showed the transformations of pearlite, bainite, and martensite, respectively. The photos by in situ observations of pearlite, bainite, and martensite are shown in (e), (j), and (o).
judgments were reasonable at the same time. On the one way, it could be seen in the martensite formation process in Fig. 11 that the martensite laths terminated at grain boundaries or laths boundaries, indicating that the dislocation pile-ups in phase boundaries poses obstacles for martensite formation. On the other way, the martensite laths nucleated in phase boundaries, which was a direct evidence for the promotion effect of the dislocations on the martensitic nucleation. It was also found from the TEM result in Fig. 9g that the substructures of the martensite laths were dislocations, which further proved the existence of dislocations in martensite laths.
that the lamella neither grew alternately along the normal direction of cementite and ferrite layers [30], nor did they grow along the lamella direction [31,32]. Hackney et al. [33] argued that the “ledge mechanism” was the formation mechanism of pearlite, but the in-situ observations supported that the formation of pearlite layer was much closer to the “precipitation” in the austenite grains in a gradual denser way, rather than the outward growth of alternate layers of cementite and ferrite. It could be seen from Fig. 4 that bainite generated at 0.03–1 °C/s. The bainite morphology transited from feathery (0.03–0.15 °C/s) to acicular (0.15–1 °C/s) with increasing cooling rate. Taking the formation of feathery bainite at triple junctions as an example, as shown in Fig. 11f–i, ferrite firstly nucleated at grain boundaries where the dislocation density was relatively high. Since the carbon content in ferrite was lower than that in supercooled austenite, the excess carbon discharged during nucleation of ferrite would diffuse into the surrounding supercooled austenite, which decreased the transformation temperatures. Considering the spread velocity (1.5 μm/s) of surface reliefs and the feathery shape of bainite (Fig. 2g), the bainite transformation process was dominated by elemental diffusion and the growth of bainite was directional, which was consistent with the diffusive theory proposed by Loginova et al. [34]. At fast cooling rate (≥0.3 °C/s), the supercooled austenite only underwent shear martensitic transformation due to slow atomic motion and insufficient time for diffusion. The transformation of martensite, illustrated in Fig. 11k–n, indicated that nucleation and growth of the martensite laths existed. The prior nucleation sites were near grain boundaries or laths boundaries, which suggested that the transformation of martensite was involved in dislocation motions. The results of Morawiec et al. [35] showed that the dislocations represented obstacles to martensite growth, and Ibarra et al. [36] proved that the dislocations were beneficial for nucleation of martensite. Therefore, the role of dislocations in martensite transformation was very complicated. The results of this study showed that these seemingly contradictory
3.5. Transformation Kinetics The in-situ observation method allowed a direct acquisition of the transformation behaviors of steels and made it possible for a direct time-temperature dependent quantitative phase amount determination. In contrast the dilatometric technique was an indirect method. An insitu quantitative analysis of phases transformation was not always possible, but in some cases when the transformation process was well recorded and processed, the phase transformation amount could be determined. In the following, as an example of phase transition kinetics determination via image processing, the bainite transformation kinetics will be described at the cooling rate of 0.3 °C/s. Fig. 12a showed the microstructures above 493 °C. The bainite formed at different temperature ranges could be painted to different colors, as shown in Fig. 12b–f. It could be seen that the phase transition amount in Fig. 12b was relatively low compared to that in Fig. 12c and d. A massive phase transition started in Fig. 12c and terminated in Fig. 12e. It showed that only a few bainite was transformed in Fig. 12f. The overlapping of Fig. 12b–f was shown in Fig. 12g. The bainite transformation volume fraction could be determined via counting pixels of different colors at different temperatures, as shown in Fig. 12h. The relationship between the bainite transformation volume fraction (fb, 0–1) and the temperature (T, °C) could be fitted by [37]: 408
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Fig. 12. Phase transition amount determination via image processing of in-situ observation results at the cooling rate of 0.3 °C/s. The temperature ranges in which the phase transformed were (a) ≥493 °C, (b) 493–461 °C, (c) 461–429 °C, (d) 429–397 °C, (e) 397–365 °C, (f) 365–333 °C, (g) the overlapping of transformed phases, and (h) transition kinetics. n
k (T − T0 ) ⎞ ⎞ ⎟ ⎟ fb = 1 − exp ⎜⎛−⎜⎛ β ⎠ ⎠ ⎝ ⎝
forgings, the cooling rates vary from part to part, therefore the precise prediction of forging microstructures and performances is very difficult. With the help of the HTCLSM technology, the microstructure evolution during the whole process of heating, forging, and subsequent cooling can be conveniently predicted. Firstly, the microstructure evolutions in heating process can be visualized on HTCLSM [22]. Then, the compression loads could be applied to obtain various strains, and the microstructures evolutions could be in-situ observed on HTCLSM. Finally, in the cooling process, the phase transformations could also be seen. Hardness, retained austenite content, and microstructures at room temperature could be predicted by the result of this investigation. In industrial applications, according to the result of this investigation, good machinability can be obtained at cooling rate lower than 0.03 °C/ s. Because lower bainite has excellent overall mechanical properties, so, in order to get a combination of high strength and high fatigue performance, cooling rate can be controlled at 0.1–1 °C/s. Technical challenges of high-temperature metallography may also be faced for HTCLSM since it is developed from the high-temperature metallography. For example, decarburization is very hard to be avoided especially in long-time in-situ heat-treatment observations, and therefore, the phase change points are relatively more accurate in fast cooling than in slow cooling. The surface reliefs generated in previous transformations can make the subsequent phase change observation
(3)
where k, β, n, and T0 denoted the frequency factor, the cooling rate (0.3 °C/s), the Avrami exponent, and the starting temperature of bainite transformation (493 °C), respectively. The frequency factor (k) and the Avrami exponent (n) were determined to be −0.00357 and 1.8946 via curve fitting. The R value in representative of the confidence level was calculated to be 0.988. It could be seen from the fitted curve in Fig. 12h that the calculated bainite transformation volume fraction agreed well with experimental value, indicating that the model was accurate in bainite transformation volume fraction. The in-situ observation method was usually adopted in high temperature austenite grain size determination, but very few researches have reported the direct application on quantitative analysis of phases transformation, and in this investigation, the bainite transformation was accurately constructed via in situ observations. 3.6. Application Prospects and Technical Challenges The 300M steel heavy forgings are often cooled in furnace, air or pit in engineering practices, and it usually takes several hours to one day to cool. Considering the differences in the shapes and thicknesses of the 409
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difficult, which remains an issue to be solved. However, in spite of these shortcomings, the HTCLSM technology is still an effective, direct, and accurate method for phase transition characterizations.
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4. Conclusions The continuous cooling transformation behaviors of 300M steel were investigated by HTCLSM, SEM, TEM, XRD, microhardness test, and dilatometry. The following conclusions could be drawn: (1) The phase transitions at various cooling rates were in-situ observed, and the cooling rate ranges of pearlite, bainite and martensite were determined to be 0.01–0.15, 0.03–1, 0.3–100 °C/s, respectively. The continuous cooling transformation diagram of 300M steel was constructed according to in-situ observations and verified by traditional methods. (2) The continuous cooling transformation diagram constructed in this investigation was compared with the results of dilatometry. The transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher, which may be due to the fact that some transformations might have taken place prior to the observed transformation points and not been observable because of a too low surface relief. (3) Pearlite was observed to form in a way that the newly formed lamella densified the lamella formed at higher temperatures. This can be explained by the carbon diffusion under the driving force of the reduction of free energy. (4) The relationship between the Vickers hardness, retained austenite content, and cooling rates was determined experimentally, and the phase transformation kinetics of martensite was formulated via insitu observations. Acknowledgements The authors appreciate the financial support by the National Natural Science Foundation of China (No. 51435007, 51705169) and the China Postdoctoral Science Foundation (No. 2017M610472). The authors would like to thank Ms. Qianqian Zhu from Advanced Manufacturing and Technology Experiment Center in Huazhong University of Science and Technology for the technical support in TEM, and Mr. Wang from Chengdu Century Mei Yang Technology Co. (http://www.ceshigo.com) for the technical support on XRD tests. The first author wants to express his gratitude to his wife, Ms. Jiao Zeng, for the companionship and support during the hard time. Appendix A. Supplementary Data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.matchar.2018.07.034. References [1] G.R. Speich, Metals Handbook Vol. 1 Properties and Selection Irons and Steels, Ultrahigh-Strength Steels Section, ASM International, Materials Park, 2005, pp. 1118–1170. [2] R. Chen, Z. Zheng, J. Li, et al., Crystals (8) (2018) 282, https://doi.org/10.3390/ cryst8070282.
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In-situ investigation of phase transformation behaviors of 300M steel in continuous cooling process
T
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Rongchuang Chen, Zhizhen Zheng , Ning Li, Jianjun Li, Fei Feng School of Materials Science and Engineering, Huazhong University of Science and Technology, and State Key Laboratory of Materials Processing and Die & Mould Technology, 430074, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Microstructure evolution Phase transition In-situ investigation 300M steel
Comprehensive understanding of continuous cooling transformation behaviors is important in heat treatment of steels, but there is still a lack of research in the phase transition behaviors of 300M steel, which would greatly hinder the microstructure controlling and finally affect the application of this material. In this work, in-situ observations by high temperature confocal laser scanning microscope (HTCLSM) was introduced by considering its superiorities in clarity, accuracy, and directness, to systemically investigate the microstructure evolutions of 300M steel under various cooling rates (0.01–100 °C/s). Pearlite, bainite, and martensite were observed to form at the cooling rate range of 0.01–0.15, 0.03–1, and 0.3–100 °C/s, respectively. A continuous cooling transformation diagram was constructed based on the in-situ observations, verified by metallography, and compared with dilatometry. Results showed that the transition temperatures by in-situ observations agreed well with the results of dilatometry, while the transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher. Finally, models accurately describing the relationships between Vickers hardness, retained austenite content, and cooling rates were established, and the phase transformation mechanisms and kinetics were analyzed. Our finding not only understands fundamentally the transformation mechanisms of different microstructures, but also provides a useful reference for practical microstructure control in heat treatment of 300M steel.
1. Introduction Steel, as a high strength and economic material, has played an extraordinary role in the development of modern industry. As a kind of ultra-high strength steel (σb ≥ 1800 MPa), medium alloyed 300M steel has been widely used in aircraft landing gear, airframe parts, and pressure vessels owing to its excellent mechanical properties [1]. However, 300M steel is facing difficulties in microstructure control in heat treatment, which seriously affects the mechanical properties and the safety of components in service. In general, the microstructures of steel could be theoretical analyzed and controlled through heat treatment, therefore the continuous cooling transformation (CCT) diagram or time-temperature-transformation (TTT) diagram becomes crucial for steels [2]. Compared with the TTT diagram, the CCT diagram has attracted more attention because it has more similarity to the cooling process in industrial applications. Although the CCT diagram of 300M steel was established by Dong et al. [3], very few in-situ observations of phase transformations of 300M steel were found, and the researches on the phase transformation behaviors of 300M steel are still needed.
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Various methods were adopted by researchers to gain a better insight into the phase transformation behaviors of steels. In conventional continuous cooling and interrupted cooling method, several groups specimen were separately cooled at constant cooling rates and quenched at different temperatures. The starting and terminating temperatures for phase transformations were determined by etching and metallography, and subsequently the CCT diagram was constructed [4]. It is worth noting that quenching might destroy the high temperature microstructures and therefore affect the accuracy of results [5]. Therefore, efforts have been made to solve this problem. One is dilatometry, in which a dilatometer recorded the thermal expansion as a function of temperatures, and then the phase transition temperatures could be determined [6,7]. Other methods such as ultrasonic techniques, high temperature X-ray diffraction (XRD), and differential scanning calorimeter analysis were also implemented to characterize the phase transformations of alloys [8–11], which significantly reduced the number of experiments. However, some factors such as thermal expansion, sound wave propagation velocity, and the width of the XRD profile were influenced by grain growth and dislocation density as well
Corresponding author. E-mail address: [email protected] (Z. Zheng).
https://doi.org/10.1016/j.matchar.2018.07.034 Received 3 May 2018; Received in revised form 14 July 2018; Accepted 26 July 2018 Available online 27 July 2018 1044-5803/ © 2018 Published by Elsevier Inc.
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as phase transitions, which inevitably introduced experimental error [12,13]. In addition, if two phases precipitated successively in cooling, it was very hard to distinguish between their phase transformation temperatures from the measured curves [6,14]. Furthermore, the relationships between temperatures and properties, such as the thermal expansion and specific heat capacity, were usually nonlinear even under no phase transformation conditions, which caused difficulty in transition temperature determinations [1]. Besides, the phase transformations were not directly observed by these methods, and the results by these methods should be further validated by metallographic observations. Accordingly, a more direct method is urgently needed. Comparing with these conventional methods, in-situ observation provided precise, direct, and continuous observations without altering the testing conditions. In general, high temperature optical microscope (HTOM), high temperature scanning electron microscope (HTSEM), and high temperature confocal laser scanning microscope (HTCLSM) are the most successful techniques for in-situ observations. But very few scanning electron microscopes could be adopted in high temperature observation of steels, which hindered the engineering applications in phase transformation researches [15]. Several in-situ observation works of phase transformations based on HTOM have been carried out, but difficulties were encountered when the phase transformation mechanisms needed to be explained in detail at higher magnifications (≥500) [16–18]. The confocal laser scanning microscope (CLSM) improved the optical imaging precision by using a spatial pinhole to block out-offocus light in image formation and has been successfully used in life sciences, pharmacology, and materials research, etc. The high temperature confocal laser scanning microscope (HTCLSM), an instrument developed for metal high-temperature metallographic observation based on CLSM, could simulate the heat treatment process of industrial conditions via accurate process controlling in protective atmospheres. The HTCLSM had its distinctive advantages in imaging clarity comparing with HTOM, and compared with HTSEM, the process control accuracy was higher, but the cost was much lower. Recently, several researches on alloys solidification, phase transformation, and grain growth have been reported [19–22] based on the HTCLSM. Although the phase transformation and grain growth of various alloys have been investigated, there is no in-situ research on the phase transformation of 300M steel. Accordingly, the HTCLSM was adopted to investigate the continuous cooling transformation behaviors of 300M steel. The detailed phase transformation behaviors at various cooling rates will be investigated. The CCT diagram of 300M steel will be established according to in-situ observation and compared with the result of dilatometry. This investigation intends to establish a detailed and systematic description of the phase transformation behaviors of 300M steel during continuous cooling.
2.2. XRF and SEM-EDS Tests The compositions were measured via an XRF apparatuses (XRF1800, Shimadzu Inc., Kyoto, Japan), and the error of element content measuring was claimed to be within ± 0.01% by the instrument supplier. The compositions in Table 1 showed that the segregation basically did not exist in the as-received ingot, although the carbon content of half radius was a little higher, which was within the allowable range of experimental error. The microstructures of as-received material and heat-treated material were observed on a scanning electron microscope (SEM, Quanta650 FEG, FEI Co., America) and the chemical compositions were quantitatively analyzed by cathodoluminescence spectroscopy (MP-32S, Horiba Co, Japan). 2.3. In-situ Observation Experiments The microstructure transformations during continuous cooling of 300M steel were in-situ observed by HTCLSM. The HTCLSM setup (VL2000DX-SVF17SP, Yonekura MFG Co., Japan), schematically illustrated in Fig. 1c, used a spatial pinhole to block out-of-focus light on a laser scanning microscope to achieve high resolution (0.14 μm) imaging. The Maximum scanning speed was 120 frames per second, ensuring that enough intermediate steps could be captured during phase transitions. In the case of quenching at 100 °C/s, the cooling time was ~11 s, and the maximum number of frames that could be captured was 1320, meaning that the HTCLSM could still record the martensitic transformation process by capturing one picture per degree centigrade decreasing. The temperature was measured using a thermocouple connected to the bottom of ceramic crucible on which the observed specimens were put. The temperature was continuously adjusted by a computer to obtain the specified temperatures and cooling rates. The temperature controlling error was claimed to be within ± 0.1 °C by the instrument supplier. During continuous cooling, the boundaries of different grains and phases at high temperatures could be visualized because of the volatilization of low melting point alloying elements at interfaces and the surface reliefs generated due to phase transformation. It was reported that the surface reliefs for bainite were inversed ‘V’ shapes and for martensite they were ‘N’ shapes [1,23]. The specimens were heated at 1 atm argon atmosphere to 1150 °C at a constant rate of 3.33 °C/s, held for 10 min, and then cooled to 25 °C at various rates ranging from 0.01, 0.03, 0.05, 0.1, 0.15, 0.3, 0.5, 1, 2, 3.33, 10, to 100 °C/s. The test specimens for HTCLSM were cut from ingot and turned to the size of 6 mm in diameter and 3 mm in height, and then mechanically polished according to standard metallographic procedure [24]. 2.4. XRD Tests
2. Materials and Experiments
The retained austenite contents of heat-treated specimens were quantitatively analyzed via an X-ray diffraction instrument (D8 Advance, Bruker Co., Germany). The specimens for XRD tests were square sheets with the size of 10 mm × 10 mm × 2 mm, and the measured surfaces were mechanically polished. The scanning speed was 1°/ min. The Co target was used in XRD experiments. The retained austenite contents were calculated via a spectrum analysis software (Topas, Bruker Co., Germany) automatically.
2.1. Materials The 300M steel ingot was in an annealed state with the diameter of 300 mm. The chemical compositions of the ingot were measured via XRay Fluorescence (XRF). In order to investigate the segregation behaviour of the ingot, specimens for XRF tests were sampled from the center, the half radius, and the outer edge of the ingot by wire-electrode cutting. Then the specimens were turned to 33 mm in diameter and 12 mm in height, and mechanically polished. The chemical compositions on different locations of the ingot were shown in Table 1. A scanning electron microscope integrated with energy dispersive spectrometer (SEM-EDS, Quanta650 FEG, FEI Co., America) was used to characterize both microstructure and chemical compositions, as shown in Fig. 1a and b. The microstructure was upper bainite. The chemical compositions measured by EDS (Fig. 1b) agreed well with the XRF result.
2.5. Microhardness Tests Micro-hardness of cooled specimens was tested on a Wilson Hardness tester (430SVD, Buehler Co., America). The test error was claimed to be within ± 0.75%, but the errors could be increased because the diagonal lengths of indentations were measured by visual inspection. The loading force was 9.8 N, and the loading time was 10 s. At least 5 points on each specimen were tested. 401
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Table 1 Chemical compositions of 300M steel (mass fraction %). Element
C
Mn
Si
Ni
Cr
V
Mo
S
Fe
Center Half radius Outer edge
0.39 0.40 0.39
0.801 0.810 0.814
2.707 2.362 2.618
1.822 1.833 1.818
0.897 0.901 0.891
0.085 0.087 0.085
0.432 0.435 0.437
0.019 0.016 0.017
Balanced Balanced Balanced
2.6. OM, SEM, and TEM Characterizations
3. Results and Discussion
Since the conventional etching and metallographic analysis were proved effective in most conditions, the cooled specimens at different cooling rates were characterized by etching and observation on OM and SEM. Specimens for OM observations were mechanically polished and etched using etchant which contained 4% HNO3 and 96% CH3CH2OH (volume fraction). An optical microscope (VHX-1000C, Keyence Co., Japan) was used for observations. Specimens were then photographed on a scanning electron microscope (SEM, Quanta650 FEG, FEI Co., America). The acceleration voltage was 10 kV. Samples for TEM tests were mechanically polished to the thickness of ~70 μm, and electro polished to obtain the thin areas for photography. The etchant for electro polishing contained 10% HClO4 and 90% CH3CH2OH. The operating voltage was 20 V, and the temperature was −20 °C. A transmission electron microscope (TECNAI G2 F30, FEI Co., America) was used in microstructure characterization. The acceleration voltage was 200 kV.
3.1. The Construction of CCT Diagram by In-situ Observations Fig. 2 showed the microstructure transformations with decreasing temperatures under various cooling rates. For example, under cooling rate of 0.01 °C/s, lamellar structures gradually formed in austenite grains with lamellar spacing decreased from 6.9 to 2.4 μm as temperature decreased from 840 (Fig. 2a) to 781 °C (Fig. 2b), and this tendency became conspicuous, such as the lamellar was almost unrecognizable on HTCLSM when the temperature further decreased to 709 °C (Fig. 2c), and the morphology of the lamellar structure remained unchanged thereafter. The interface marked in Fig. 2a–c, also named as growth ledge in literature [25], did not migrate during cooling. Under cooling rate of 0.15 °C/s, some part of the grains became darker from 753 (Fig. 2d) to 702 °C (Fig. 2e), which suggested the reflectivity changing owing to the microstructure transformations of the grains. It was also observed in Fig. 2f that upper bainite was generated when temperature decreased from 492 to 345 °C while the appearance of some grains remained unchanged throughout the cooling. Under cooling rate of 0.3 °C/s, no microstructure transformation could be seen until the temperature drops to 493 °C (Fig. 2g), which located in the temperature range of bainite transformation, surface reliefs were observed to form at an average speed of about 1.5 μm/s across grains and at grain boundaries. Fig. 2h and i showed the detailed transformation progress. Under the cooling rate of 3.33 °C/s, the microstructure transformation only occurred below 295 °C (Fig. 2j), wherein lath
2.7. Thermal Expansion Tests To compare with the phase transition points determined by HTCLSM observations, the thermal expansions were recorded by dilatometry. Specimens were turned to cylinders with the size of 8 mm in diameter and 12 mm in height. Thermal expansions were measured on a thermal compression machine (Gleeble 3500, Dynamic Systems Inc., America) at cooling rates of 0.15 and 0.3 °C/s, respectively.
Fig. 1. As-received material and experiment schematic. (a) SEM graph showed the microstructure was bainite; (b) EDS spectrum of as-received 300M steel; (c) schematic illustrations of the HTCLSM setup. 402
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Fig. 2. Microstructure transformations at cooling rates of (a)–(c) 0.01 °C/s; (d)–(f) 0.15 °C/s; (g)–(i) 0.3 °C/s; (j)–(l) 3.33 °C/s by in-situ observations on HTCLSM. “Lath A–F” in Fig. 2j–l denoted martensite laths. “GB” meant grain boundary. “UB” was upper bainite. The corresponding temperatures were marked below the photos. (×100).
divided into 7 regions, the microstructures of each were shown. The upper right region (A) showed the equiaxed austenite grains before transformations. At slow cooling rates, pearlite lamellas and proeutectoid ferrite (P + F) were the dominant structures. As the cooling rate increased, the microstructure transited from mixed microstructures of pearlite and bainite (P + B), bainite and martensite (B + M), to martensite (M). It should be noticed that the incomplete transformation was a common characteristic of bainite and martensite transformations, so there was retained austenite in the transition product as well, which should be further investigated. The regions marked ‘A → B’ and ‘A → M’ showed the intermediate state of bainite and martensite transformations. The pearlite, bainite, and martensite formed at the cooling rate range of 0.01–0.15, 0.03–1, and 0.3–100 °C/s, respectively. The transformation starting temperatures were within the range of 753–840,
martensite appeared very fast across grains and grew more and more as temperature decreased, as shown in Fig. 2k and l, respectively. It was interesting that the newly formed martensite laths were connected end to end, and almost all laths terminated at grain boundaries (marked in Fig. 2j) or where they met (Fig. 2k). The zoomed views of the pearlite formation process at the cooling rate of 0.01 °C/s were shown in Fig. 3. In the starting stage of pearlite formation, lamellas with relative big spaces were formed (Fig. 3a) in some austenite grains, and the zoomed views in Fig. 3b showed the oval nuclei between pearlite lamellas. With decreasing temperature, it could be seen from Fig. 3c and d that the newly formed pearlite lamellas densified the previously formed pearlite lamellas. To further investigate the effect of cooling rates on microstructure transformations, the HTCLSM results were summarized in Fig. 3. It was 403
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Fig. 3. Zoomed views of the pearlite formation process at the cooling rate of 0.01 °C/s and photographed at (a) 832 °C, (b) 832 °C, (c) 756 °C, and (d) 730 °C.
444–496, and 295–297 °C. At cooling rate of 0.03–0.15 (pearlite and bainite transformation) and 0.3–1 °C/s (bainite and martensite transformation), two transformations occurred successively. The XRD profiles of cooled specimens at various cooling rates were shown in Fig. 5a, according to which the retained austenite contents could be calculated, shown in Fig. 5b. The retained austenite was marked out using face centered γ-Fe, and the α-Fe represented body centered ferrite or martensite depending on the cooling rates. As could be seen, with increasing cooling rate, the intensity of γ-Fe peaks decreased gradually, indicating the retained austenite content decreased. Sherman et al. [26] reported an inverse liner relationship between the retained austenite contents and cooling rate of a martensitic steel in the cooling rate range of 25–560 °C/s. The result of Li et al. [27] also showed that the retained austenite content of 0.19C-1.5Si-1.6Mn steel decreased with increasing cooling rate from 0.05 to 10 °C/s. In the present investigation, in a wider range of cooling rates, the retained austenite content (Cra) was found to decrease with cooling rate (v), and the relationship between the retained austenite contents and the logarithm of cooling rate could be fitted using a second order polynomial:
Cra =
2 0.009496⋅log10 v
− 0.04329⋅log10 v + 0.04563
Fig. 4. Relationships between microstructure transformations and cooling rates. The letters “A”, “P”, “B” and “M” denoted austenite, pearlite, bainite, and martensite, respectively.
(1)
The symbols and line in Fig. 5b showed the experimental and fitted values. The R value in representative of the confidence level was calculated to be 0.937, indicating that the Eq. (1) was moderately precise in calculating the retained austenite content. The influence of cooling rates on retained austenite content could be explained by Fig. 4. At slower cooling rates, the transformations of γ-Fe to α-Fe were slower as well, resulting in a relative longer time to complete the transformations, and a higher content of retained austenite was obtained. Meanwhile, at slower cooling rates, the strains generated in phase transformations, which was beneficial for phase transformation according to Zhang [28], were more easily to be eliminated by recovery. Therefore, the retained austenite content was higher in the slow cooling. The micro-hardness of cooled specimens was tested, as shown in
Fig. 6. Basically, the relationships between the Vickers hardness and the logarithm of cooling rates followed the S-curve. The hardness increased with increasing cooling rates until at 3.33 °C/s the hardness became stable at HV 665 ± 15. It was confirmed by HTCLSM results that the main transformation product was pearlite, bainite, and martensite at the cooling rate range of 0.01–0.15, 0.03–1, 0.3–100 °C/s, respectively. Therefore, the bainite seemed to be an intermediate of pearlite and martensite in terms of hardness. The relationships between cooling rates (v) and Vickers hardness (Hv) could be fitted as:
Hv = 667.92 −
404
286.65 1 + exp(2.3013⋅log10 v + 1.4117)
(2)
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Fig. 5. Retained austenite content measured by XRD. (a) XRD profiles; (b) retained austenite content in percentage. Table 2 Phase transition points. The “M”, “B”, and “P” denoted martensite, bainite, and pearlite. The subscript s and f denoted phase transition starting and finishing temperatures, respectively. The symbol “-” denoted the absence of the corresponding transformations. v (°C/s)
0.01 0.03 0.05 0.1 0.15 0.3 0.5 1 2 3.33 10 100
Fig. 6. Hardness of specimens cooled at different cooling rates.
Phase transition temperatures (°C) Ms
Mf
Bs
Bf
Ps
Pf
– – – – – 296 295 295 295 295 299 297
– – – – – 160 160 164 169 160 160 162
– 496 492 496 492 493 460 444 – – – –
– 369 359 352 345 329 311 299 – – – –
840 821 800 770 753 – – – – – – –
709 708 706 708 702 – – – – – – –
began not to change. The transformation starting and finishing temperatures of martensite almost kept constant at 295 and 160 °C when the cooling rate varied from 100 to 0.3 °C/s, but the martensite transformation disappeared during cooling at a rate higher than 0.3 °C/s, so at that point, the Ms and Mf points were connected with a dotted line, indicating a termination of martensite transformation. With increasing cooling rates from 0.03 to 1 °C/s, the transformation starting and finishing temperatures of bainite decreased from 496 to 444 and from 369 to 299 °C, respectively. The transformation starting temperatures of pearlite also showed a significant decrease from 840 to 753 °C as the cooling rates increased from 0.01 to 0.15 °C/s while the finishing temperature almost kept constant at 706 °C by the HTCLSM result. It could be explained that for diffusion type transformations, for example, the bainite and pearlite transformations, the transformation temperatures would be lower at fast cooling rates due to insufficient diffusions. Besides, with decreasing cooling rates, the lower-bainite transformation and the upper-bainite transformation occurred sequentially, and the upper-bainite transformation temperatures were usually higher than the lower-bainite transformation temperatures.
Fig. 7. Continuous cooling transformation diagram of 300M steel.
The R value is 0.962, which meant the model could calculate the Vickers hardness of 300M steel with high precision in a wide cooling rate range from 0.01 to 100 °C/s. According to the in-situ observations, the CCT diagram was constructed and shown in Fig. 7, and the corresponding phase transition points were shown in Table 2. The starting points were estimated to be the first occurrence of newly generated microstructures, and the termination points were the temperatures at which the microstructure
3.2. Verification by Traditional Methods Since the conventional etching and metallographic analysis were proved effective, the cooled specimens at different cooling rates were 405
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Fig. 8. Etching results of 300M steel austenitizing at 1150 °C for 10 min and cooled at (a) 0.01; (b) 0.15; (c) 0.3; (d) 3.33 °C/s. PF denoted proeutectoid ferrite. BF meant bainite ferrite. (×200).
3.3. Comparison With the Results of Dilatometry
characterized by etching and then observed on OM. Fig. 8 showed the etching results of cooled specimens at various cooling rates. It showed that at the cooling rate of 0.01 °C/s (Fig. 8a), the equilibrium microstructures containing proeutectoid ferrite and pearlite with lamellar structures were formed. At 0.15 °C/s (Fig. 8b), the mixed microstructures of bainite ferrite and lamellar pearlite were formed. At 0.3 °C/s (Fig. 8c), lower bainite was dominant and some martensite laths were observed as well. At cooling rates of 3.33 °C/s, lath martensite was the main microstructure, as shown in Fig. 8d. The microstructures characterized by traditional methods agreed well with the result by HTCLSM. In order to scrutinize the surface morphology and substructure and get a better understanding of transformation mechanisms of different microstructures, cooled specimens obtained at various cooling rates were characterized by SEM and TEM and shown in Fig. 9. At 0.05 °C/s, mixed microstructures of pearlite and bainite were observed (Fig. 9a). The shape of pearlite was irregular lamellae, which indicated spheroidization of the cementite (Fig. 9b). It also showed that acicular bainites were generated and ferrites were distributed between bainites (Fig. 9c). At cooling rate of 0.15 °C/s, feathery bainites formed at grain boundaries (Fig. 9d). It was found that the most frequent angle between bainite leaves and grain boundaries were 43°. When cooled at 100 °C/s, martensite laths with the angle of 57–72° were generated (Fig. 9e). It could be seen from Fig. 9f that the martensite laths were blocked by the laths with angle of ~65°. The substructures of martensite lath were dislocations, as shown in the dark areas in Fig. 9g that the dislocation piled up in a martensite lath. The microstructures characterized by SEM and TEM agreed well with the result by HTCLSM.
In order to inspect the reliability of the CCT diagram established in by in-situ observations, the results were compared with the results of dilatometry. Thermal expansions were measured at the cooling rate of 0.15 and 0.3 °C/s and compared with the HTCLSM results, as shown in Fig. 10, where the phase transition temperatures obtained by HTCLSM observations were marked as well. Although the phase transition points could be observed by the thermal expansion method, it was very difficult to obtain accurate phase change points by dilatometry method especially in the stage when martensite and bainite were transformed successively. The transition points by in-situ observations agreed well with the results of dilatometry. But the transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher. It could be explained that some transformations might have taken place prior to the observed transformation points and not been observable because of a too low surface relief. The same reason was applicable to the explanation of the termination points of phase transitions. Although it may be less rigorous to declare that the result of in-situ observation was more accurate than the result of dilatometry, the in-situ observations by HTCLSM does have an obvious advantage in terms of directness and convenience.
3.4. Phase Transition Mechanisms Uneven carbon distribution existed in austenite grains before the transformation, which can be proved by the result in Fig. 2a that the
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Fig. 9. SEM and TEM results of 300M steel showing the various microstructures at different cooling rates: (a) 0.05 °C/s, P + B, SEM, ×750; (b) 0.05 °C/s, P, SEM, ×4000; (c) 0.05 °C/s, B, SEM, ×2000; (d) 0.15 °C/s, P + B, SEM, ×750; (e) 100 °C/s, M, SEM, ×2000; (f) 100 °C/s, M, TEM, ×3800, dark field image; (g) 100 °C/s, M, TEM, ×100,000, bright field image.
uneven carbon distribution was aggravated by the long-range carbon diffusion at a slow cooling rate, as shown in Fig. 8a that microstructures consisted of proeutectoid ferrite and spheroidal pearlite were formed at 0.01 °C/s. At moderate cooling rate, the proeutectoid ferrite was not formed because the long-range diffusion of carbon could not be completed due to lack of time. Therefore, the carbon was retained in ferrites, and some of the carbon precipitated from ferrite, forming cementite. This was proved by the results in Fig. 9c and d that the cementite distributed among bainite ferrite. The transition processes of pearlite, bainite, and martensite were illustrated in Fig. 11 according to the in-situ observations. As shown in Fig. 11a–d, at a slow cooling rate of 0.01–0.15 °C/s, as the solubility of carbon in supercooled austenite decreased with increasing supercooling degree, equidistant, parallel, and necklaced cementite nucleated in austenite grains. Then, the coarse cementite lamella with several microns spacing were initially formed, consuming carbon in super cooled austenite. As a result, the carbon content of the regions far from the cementite layers was relatively high comparing with the regions adjacent to the cementite layers, so the cementite nucleated and grew, leading to a reduction in lamellar spacing. Fig. 11e showed that cementite nucleated between cementite layers. It should be noticed that the pearlite lamella formed in a gradually denser way, as shown in Fig. 11a–d, and the newly formed lamella densified the lamella formed at higher temperatures. The so called “growth ledge” [25] in Fig. 2a–c did not migrate during the formation of pearlite lamella, which meant
Fig. 10. Comparison of the phase transition points obtained by in-situ observations and by dilatometry during cooling at 0.15 and 0.3 °C/s.
austenite grains did not transform with simultaneously. The higher carbon content, the bigger the resistance of phase transformation, and therefore the lower the phase transition temperatures would be. It was reported by Vermeulen et al. [29] that the Ms would decrease by about 275–375 °C for every 1% of equivalent carbon content decreasing. The 407
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Fig. 11. Illustration of the pearlite, bainite, and martensite formation. Figures in (a)–(d), (f)–(i), and (k)–(n) showed the transformations of pearlite, bainite, and martensite, respectively. The photos by in situ observations of pearlite, bainite, and martensite are shown in (e), (j), and (o).
judgments were reasonable at the same time. On the one way, it could be seen in the martensite formation process in Fig. 11 that the martensite laths terminated at grain boundaries or laths boundaries, indicating that the dislocation pile-ups in phase boundaries poses obstacles for martensite formation. On the other way, the martensite laths nucleated in phase boundaries, which was a direct evidence for the promotion effect of the dislocations on the martensitic nucleation. It was also found from the TEM result in Fig. 9g that the substructures of the martensite laths were dislocations, which further proved the existence of dislocations in martensite laths.
that the lamella neither grew alternately along the normal direction of cementite and ferrite layers [30], nor did they grow along the lamella direction [31,32]. Hackney et al. [33] argued that the “ledge mechanism” was the formation mechanism of pearlite, but the in-situ observations supported that the formation of pearlite layer was much closer to the “precipitation” in the austenite grains in a gradual denser way, rather than the outward growth of alternate layers of cementite and ferrite. It could be seen from Fig. 4 that bainite generated at 0.03–1 °C/s. The bainite morphology transited from feathery (0.03–0.15 °C/s) to acicular (0.15–1 °C/s) with increasing cooling rate. Taking the formation of feathery bainite at triple junctions as an example, as shown in Fig. 11f–i, ferrite firstly nucleated at grain boundaries where the dislocation density was relatively high. Since the carbon content in ferrite was lower than that in supercooled austenite, the excess carbon discharged during nucleation of ferrite would diffuse into the surrounding supercooled austenite, which decreased the transformation temperatures. Considering the spread velocity (1.5 μm/s) of surface reliefs and the feathery shape of bainite (Fig. 2g), the bainite transformation process was dominated by elemental diffusion and the growth of bainite was directional, which was consistent with the diffusive theory proposed by Loginova et al. [34]. At fast cooling rate (≥0.3 °C/s), the supercooled austenite only underwent shear martensitic transformation due to slow atomic motion and insufficient time for diffusion. The transformation of martensite, illustrated in Fig. 11k–n, indicated that nucleation and growth of the martensite laths existed. The prior nucleation sites were near grain boundaries or laths boundaries, which suggested that the transformation of martensite was involved in dislocation motions. The results of Morawiec et al. [35] showed that the dislocations represented obstacles to martensite growth, and Ibarra et al. [36] proved that the dislocations were beneficial for nucleation of martensite. Therefore, the role of dislocations in martensite transformation was very complicated. The results of this study showed that these seemingly contradictory
3.5. Transformation Kinetics The in-situ observation method allowed a direct acquisition of the transformation behaviors of steels and made it possible for a direct time-temperature dependent quantitative phase amount determination. In contrast the dilatometric technique was an indirect method. An insitu quantitative analysis of phases transformation was not always possible, but in some cases when the transformation process was well recorded and processed, the phase transformation amount could be determined. In the following, as an example of phase transition kinetics determination via image processing, the bainite transformation kinetics will be described at the cooling rate of 0.3 °C/s. Fig. 12a showed the microstructures above 493 °C. The bainite formed at different temperature ranges could be painted to different colors, as shown in Fig. 12b–f. It could be seen that the phase transition amount in Fig. 12b was relatively low compared to that in Fig. 12c and d. A massive phase transition started in Fig. 12c and terminated in Fig. 12e. It showed that only a few bainite was transformed in Fig. 12f. The overlapping of Fig. 12b–f was shown in Fig. 12g. The bainite transformation volume fraction could be determined via counting pixels of different colors at different temperatures, as shown in Fig. 12h. The relationship between the bainite transformation volume fraction (fb, 0–1) and the temperature (T, °C) could be fitted by [37]: 408
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Fig. 12. Phase transition amount determination via image processing of in-situ observation results at the cooling rate of 0.3 °C/s. The temperature ranges in which the phase transformed were (a) ≥493 °C, (b) 493–461 °C, (c) 461–429 °C, (d) 429–397 °C, (e) 397–365 °C, (f) 365–333 °C, (g) the overlapping of transformed phases, and (h) transition kinetics. n
k (T − T0 ) ⎞ ⎞ ⎟ ⎟ fb = 1 − exp ⎜⎛−⎜⎛ β ⎠ ⎠ ⎝ ⎝
forgings, the cooling rates vary from part to part, therefore the precise prediction of forging microstructures and performances is very difficult. With the help of the HTCLSM technology, the microstructure evolution during the whole process of heating, forging, and subsequent cooling can be conveniently predicted. Firstly, the microstructure evolutions in heating process can be visualized on HTCLSM [22]. Then, the compression loads could be applied to obtain various strains, and the microstructures evolutions could be in-situ observed on HTCLSM. Finally, in the cooling process, the phase transformations could also be seen. Hardness, retained austenite content, and microstructures at room temperature could be predicted by the result of this investigation. In industrial applications, according to the result of this investigation, good machinability can be obtained at cooling rate lower than 0.03 °C/ s. Because lower bainite has excellent overall mechanical properties, so, in order to get a combination of high strength and high fatigue performance, cooling rate can be controlled at 0.1–1 °C/s. Technical challenges of high-temperature metallography may also be faced for HTCLSM since it is developed from the high-temperature metallography. For example, decarburization is very hard to be avoided especially in long-time in-situ heat-treatment observations, and therefore, the phase change points are relatively more accurate in fast cooling than in slow cooling. The surface reliefs generated in previous transformations can make the subsequent phase change observation
(3)
where k, β, n, and T0 denoted the frequency factor, the cooling rate (0.3 °C/s), the Avrami exponent, and the starting temperature of bainite transformation (493 °C), respectively. The frequency factor (k) and the Avrami exponent (n) were determined to be −0.00357 and 1.8946 via curve fitting. The R value in representative of the confidence level was calculated to be 0.988. It could be seen from the fitted curve in Fig. 12h that the calculated bainite transformation volume fraction agreed well with experimental value, indicating that the model was accurate in bainite transformation volume fraction. The in-situ observation method was usually adopted in high temperature austenite grain size determination, but very few researches have reported the direct application on quantitative analysis of phases transformation, and in this investigation, the bainite transformation was accurately constructed via in situ observations. 3.6. Application Prospects and Technical Challenges The 300M steel heavy forgings are often cooled in furnace, air or pit in engineering practices, and it usually takes several hours to one day to cool. Considering the differences in the shapes and thicknesses of the 409
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difficult, which remains an issue to be solved. However, in spite of these shortcomings, the HTCLSM technology is still an effective, direct, and accurate method for phase transition characterizations.
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4. Conclusions The continuous cooling transformation behaviors of 300M steel were investigated by HTCLSM, SEM, TEM, XRD, microhardness test, and dilatometry. The following conclusions could be drawn: (1) The phase transitions at various cooling rates were in-situ observed, and the cooling rate ranges of pearlite, bainite and martensite were determined to be 0.01–0.15, 0.03–1, 0.3–100 °C/s, respectively. The continuous cooling transformation diagram of 300M steel was constructed according to in-situ observations and verified by traditional methods. (2) The continuous cooling transformation diagram constructed in this investigation was compared with the results of dilatometry. The transition starting temperatures by in situ observation were lower and the phase transition termination temperatures were higher, which may be due to the fact that some transformations might have taken place prior to the observed transformation points and not been observable because of a too low surface relief. (3) Pearlite was observed to form in a way that the newly formed lamella densified the lamella formed at higher temperatures. This can be explained by the carbon diffusion under the driving force of the reduction of free energy. (4) The relationship between the Vickers hardness, retained austenite content, and cooling rates was determined experimentally, and the phase transformation kinetics of martensite was formulated via insitu observations. Acknowledgements The authors appreciate the financial support by the National Natural Science Foundation of China (No. 51435007, 51705169) and the China Postdoctoral Science Foundation (No. 2017M610472). The authors would like to thank Ms. Qianqian Zhu from Advanced Manufacturing and Technology Experiment Center in Huazhong University of Science and Technology for the technical support in TEM, and Mr. Wang from Chengdu Century Mei Yang Technology Co. (http://www.ceshigo.com) for the technical support on XRD tests. The first author wants to express his gratitude to his wife, Ms. Jiao Zeng, for the companionship and support during the hard time. Appendix A. Supplementary Data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.matchar.2018.07.034. References [1] G.R. Speich, Metals Handbook Vol. 1 Properties and Selection Irons and Steels, Ultrahigh-Strength Steels Section, ASM International, Materials Park, 2005, pp. 1118–1170. [2] R. Chen, Z. Zheng, J. Li, et al., Crystals (8) (2018) 282, https://doi.org/10.3390/ cryst8070282.
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